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Grain growth acceleration during high temperature deformation in high purity alumina
Materials Science and Engineering, A 149 ( 1991 ) 59-64
59
Grain growth acceleration during high temperature deformation in high purity alumina Yu-ichi Yoshizawa and Taketo Sakuma Department of Materials Science, Faculty of Engineering, The Universityof Tokyo, 7-3-I Hongo, Tokyo 113 (Japan) (Received July 9, 1991 )
Abstract The accelerated grain growth during high temperature deformation in high purity alumina was investigated. The high purity alumina with an initial grain size about 0.6/~m was characterized by a high work-hardening rate after initial yielding at temperatures of 1250-1500 °C. The work hardening was related to the grain growth during deformation. The grain growth was accelerated by deformation which was analyzed to be of strain-enhanced type, because the grain size increase could be described as a function of strain but was not sensitive to temperature. The strain-enhanced grain growth in high purity AI203 and MgO-doped A1203 was discussed in comparison with the data for a tetragonal zirconia polycrystal reported previously.
1. Introduction Superplasticity in fine-grained ceramics, particularly in tetragonal zirconia polycrystals (TZPs), has been studied extensively by many workers [1-10]. It is essential to fabricate ceramics with a grain size below about 1/~m for realizing superplasticity in ceramics [1]. In the case of alumina (AI203), however, superplastic deformation has not been attained even for finegrained samples, because it is not easy to obtain steady state deformation owing to extensive grain growth and cavitation during high temperature deformation [11-16]. Such a microstructure change during deformation makes it difficult to understand the deformation mechanism of this material. In this study, the grain growth during high temperature deformation was examined in high purity A1203 and magnesia (MgO)doped AI203 in detail.
2. Experimental procedures High purity AI203 powders with an average diameter of 0.1 /xm supplied by Taimei Chemical Co. Ltd. (TM-DAR) and high purity MgO powders of 99.97% purity with about 17 nm diameter supplied by Ube Co. Ltd. were used as starting materials. Two kinds of material were prepared from these starting materials. One was pure A I 2 0 3 and the other was A I 2 0 3 doped with 0.1 wt.% MgO. AI203 powders were stirred with ethanol and followed by drying and passing through a 0921-5093/91/$3.50
60 mesh sieve. For fabricating MgO-doped A1203, the A1203 and MgO powders were mixed in a ball mill in ethanol together with high purity A120 3 balls for 24 h and then dried and sieved. These sieved powders were pressed into bars in a cemented carbide die under a pressure of 33 MPa and cold isostatically pressed under 100 MPa in a rubber tube. Sintering was done at 1300 °C for 2 h in air. The bulk density of the assintered and deformed specimens was measured by the Archimedes technique. Specimens of size about 7 m m x 7 m m x 7 mm were cut from the sintered bodies and ground with a diamond wheel. A high temperature mechanical test was done by uniaxial compression in a temperature range between 1250 and 1500 °C in air using an Instron-type Shimadzu AG5000C mechanical testing machine equipped with a high temperature furnace. The test was performed in a strain rate range between 1 × 10- 5 and 5 x 10- 4 s- 1. The compression test was started after holding at the test temperature for 10 rain. Small specimens, which were not deformed, were placed beside the compression test specimens to measure the static grain growth. The Pt-(Pt-Rh) thermocouples were attached to both specimens. The strain rate sensitivity exponent m was measured by the strain rate change method. Deformed and statically annealed specimens were cut and polished and were then thermally etched at 1200 °C for 4 h in air. They were coated with a gold film using an ion-sputtering machine in order to carry out scanning electron microscopy observations. The grain size was measured by the linear intercept method using © 1991 --Elsevier Sequoia, Lausanne
60
Y. Yoshizawa and T. Sakuma
/
Graingrowth acceleration during high T deformation in Al_,O~
photographs taken with a JEOL JSM-5200 scanning electron microscope. Metallographic examinations of thin foils prepared by ion milling were made with a JEOL JEM-2000FX transmission electron microscope operated at 200 kV.
3. Results
The density of the as-sintered state was about 98.5% of the theoretical value for pure and MgO-doped A1203, and the grain sizes of the two materials were 0.85/~m and 0.66/~m respectively. Figure 1 shows the transmission electron micrograph of MgO-doped A1203 sintered at 1300 °C for 2 h. The grains are nearly equiaxed and only a small number of pores is observed. No second phases such as glass phase or spinel were found, even at grain boundary triple junctions. The stress-strain curves for pure and MgO-doped A1203 deformed at 1300 °C for an initial strain rate of 1.3 x 10 -4 S- l are shown in Fig. 2. For comparison, the stress-strain curve of a TZP with an initial grain size of about 0.3/~m under the same deformation conditions is shown as a broken curve. The initial yield stresses (the 0.2% proof stress) of pure AI203 and MgO-doped A1203 are not very different from that of TZP at this temperature, but the work-hardening behaviours of A I 2 0 3 with and without MgO are significantly different from that of TZE A marked stress increase after initial yielding occurs in Al203 in contrast with TZP. The MgO addition results in a decrease in the work-hardening rate of A I 2 0 3. Figure 3 shows the stress-strain curves of MgOdoped Al203 at various temperatures. The compression test was done up to a maximum load of 20 kN.
The flow stress is largely dependent on the deformation temperature. The MgO-doped A1203 exhibits work hardening at temperatures between 1250 and 1500 °C. On the contrary, pure A1203 showed a larger work hardening than MgO-doped A I 2 0 3 as mentioned before. This large work hardening in pure AI203 often resulted in crack initiation, which is accompanied by a stress drop. Figure 4 is the transmission electron micrograph of deformed MgO-doped AI203. Grains remained equiaxed and dislocations are not observed. However, the grain size in the deformed sample is a little larger than in the as-sintered sample (Fig. 1). This microstructure was characteristic of fine-grained samples deformed at temperatures of 1250-1350 °C. In these samples, the density increased after deformation, e.g. MgO-doped AI203 with an initial density of 98.5% had densities of 99.7% and 99.9% after 30% strain at 1300 °C and 1350 °C respectively. Deformation above 1400 °C resulted in a different microstructure. Cavities were generally formed and the density
200
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..............
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Fig. 2. Stress-strain curves of pure and MgO-doped AI203 at 1300 °C: - - - , stress-strain curve for TZP under the same conditions.
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Fig. 3. Stress-strain curves of MgO-doped AI203 at various temperatures.
Y. Yoshizawa and 7'. Sakuma
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Grain growth acceleration during high T de[onnation in A 120 ~
61
Fig. 4. Transmission electron micrograph of MgO-doped AI203 31% deformed at 1300 °C. decreased after deformation above 1400 °C. In some coarse grains of size larger than about 2/~m, dislocations were introduced. Microstructural evidence indicates that it is not easy to understand the high temperature deformation mechanism in the present materials. In this study, grain growth behaviour was investigated for a grain size of less than 2/~m. Figures 5(a), 5(b) and 5(c) are the scanning electron micrographs of as-sintered, annealed and deformed MgO-doped A1203 respectively. The annealing and deformation temperature was 1300°C. In the assintered state, the average grain size was 0.66 tim. The grain sizes of both the annealed and the deformed samples are larger than that of the as-sintered sample. However, as clearly seen in Figs. 5(b) and 5(c), the deformed sample in Fig. 5(c) has a larger grain size than the annealed sample in Fig. 5(b). The average grain sizes were 0.67 ,um and 0.95 ktm for annealed and deformed samples respectively. This means that the grain growth is accelerated during high temperature deformation. The grain size change in MgO-doped AI203 during annealing and deformation at 1300 and 1350 °C is shown in Fig. 6. The terms static and dynamic are used for the grain growth during annealing and deformation in this figure. At 1300 °C, the static grain growth is almost negligible, but the grain size clearly increases with increasing deformation time. The grain growth takes place more rapidly at 1350 °C. At this temperature, the grain growth is also accelerated by deformation. 4. Discussion
It is well known that the grain growth during superplastic deformation in metals takes place faster than
Fig. 5. Scanning electron micrographs of MgO-doped AI203 (a) in the as-sintered state, (b) statically annealed and (c) 31% deformed at 1300 °C. The microstructure of (b) is obtained for a sample placed besides the deformed sample whose microstructure is (c). The time to keep at 1300 °C in the annealed sample is exactly the same as the loaded time in the deformed sample.
during static annealing [17-22]. The accelerated grain growth caused by high temperature deformation in superplastic metals is described as a function of plastic strain and is called the "strain-enhanced grain growth" [17]. The grain growth acceleration during superplastic or high temperature deformation has also been reported in TZP and AI20 ~ [2, 8, 9, 12-16]. Nieh and Wadsworth [8, 9] have indicated that the cube d 3 - d . ~ of grain size increment was proportional to the defor-
62
Y. Yoshizawa and 7". Sakuma
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Graingrowth acceleration during high Tdeformation in AI,_O.~
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Fig. 6. Grain size change during annealing and deformation at 1300 and 1350 °C for MgO-doped A1203: o, D, data for annealed samples; e,-, data for the deformed samples at an initial strain rate of 1.2 x 10 -4 S- t. The horizontal axis shows the time for annealing or deformation. mation time in TZP, where do is the initial grain size; Chen and Xue [2] have shown that In(d/do) changes with strain in Ce-TZE They have also reported that ( d - d o ) / d o is expressed as a function of deformation time in high purity A1203 [16]. Thus various plots have been proposed to analyse the accelerated grain growth in ceramics, but no theoretical background has been established in these plots. The present data in Fig. 6 clearly show that the dynamic grain growth includes both static growth and accelerated growth due to deformation. The contribution of static growth becomes more important at higher deformation temperatures. It seems therefore reasonable that the grain size difference d - d a between deformed and annealed samples should be discussed rather than d - d 0 for analysing the accelerated grain growth. The present authors have recently reported that the value of ( d - d a ) / d a increases almost linearly with strain and is not sensitive to deformation temperature in TZP [10, 23]. However, the magnitude of strainenhanced grain growth rate in TZP was about one order of magnitude lower than in most superplastic metals [10]. For examining whether or not the grain growth acceleration in high purity Al203 is expressed as a function of plastic strain, the normalized grain size increment ( d - d a ) / d a is plotted against strain in Fig. 7. The data on ( d - d a ) / d a for both pure and MgO-doped Al203 can be expressed by straight lines passing through the origin. The slopes of the lines are slightly different for pure and MgO-doped AI203 but are not sensitive to temperature. The slopes are estimated to be 2.4 and 1.5 for pure and MgO-doped Al203 respectively. These values are in good agreement with the reported value of 2 for superplastic Sn-Bi alloy
08.;
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0.,
0.a
0.8
,0
Fig. 7. The relationship between the normalized grain size difference and strain for pure and MgO-doped A I 2 0 3 . The strain rate is in the range between 1.2 x 10 -5 and 5 x 10 -4 s - 1.
[19], but about 20 times larger than the value for TZP of about 0.1 [10]. Therefore it seems valid that the enhancement of grain growth caused by deformation in A1203 is regarded as strain enhanced. In comparison with superplastic metals, the experimental data on the strain-enhanced grain growth in ceramics are very limited, so that it is not possible to give a firm interpretation for the accelerated grain growth. However, for elucidating the mechanism of strain-enhanced grain growth, it may be valuable to examine whether or not the data for AI203 can be treated by analogy to those for superplastic metals. Wilkinson and Cficeres [21] have clarified a general rule on the strain-enhanced grain growth in superplastic metals. They have found that the log-log plot of grain growth rate normalized by initial grain size against strain rate is expressed by a sigmoidal relationship for various superplastic metals and is divided into three regions. In particular, the data fall in a narrow band and are expressed as a linear relationship in the intermediate strain rate range. Figure 8 is such a plot for the present alloys. In Fig. 8, the dotted zones are the data for various superplastic metals [21] and TZP obtained by the present authors [23]. In the strain rate range in Fig. 8, the data on both high purity A1203 and MgO-doped A120 3 are represented by straight lines whose slopes are close to unity as are those for superplastic metals and TZP. However, the absolute value of growth rate in alumina is much higher than that in TZP and rather agrees with that in superplastic metals.* *In the plot of Wilkinson and Cficeres, the strain-enhanced grain growth was not estimated from d-da as in the present paper but from d - d0. If the data on superplastic metals are replotted as d - d a, the conclusion may be slightly different.
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deformed mainly by grain boundary sliding at the temperatures and strain rates examined. Then, it is rather realistic to assume that the parameters b, the average migration distance per unit sliding, are different in different materials and that b is much smaller in TZP than in A1203 and metals. It has been shown that b consists of two parameters:
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10-s
i
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i
i
i
I iiii[
i
i
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~ / s -1 Fig. 8. The relation between normalized strain-enhanced grain growth rate and strain rate. The data for metals are those compiled by Wilkinson and Cficeres [21 ], and those for TZP are the reported values obtained by present authors. The strain-enhanced grain growth rate in A1203 can be described by the following equation:
[t,. = abdg
(1)
where a is the fraction of overall strain due to grain boundary sliding, b is the average migration distance per unit sliding and d is the grain size. Equation (1) is for an intermediate strain rate range in superplastic metals, and the theoretical basis has been given by Wilkinson and Cficeres [21 ]. This equation predicts the proportionality between strain-enhanced grain growth rate and strain rate, and a linear relationship with a slope of unity in the log dr vs. log g plot being consistent with the data in Fig. 8. This may support the fact that the strain-enhanced grain growth in A1203 takes place by a similar mechanism to that in superplastic metals. Next, the difference of the proportional constant a b d in materials is discussed. Of the three parameters, the grain sizes d themselves are different in materials. A typical grain size for superplastic metals is several micrometres and that of TZP is 0.3-0.5 #m, while the grain size of the present AI203is between 0.6 and 1.5 #m. Thus it is better to discuss the enhanced grain growth by a normalized form:
d~, abg d
(2)
The proportional constant a b in eqn. (2) is estimated to be 2.5 for pure A1203, 1.4 for MgO-doped A1203, about 0.13 for TZP, and around 1 for superplastic metals. It is unrealistic to assume that the difference is caused by a, the fraction of overall strain due to sliding, because all these materials are believed to be
(3)
for grain growth enhanced by grain boundary sliding [23], where c is the depth of deformation zone produced by unit sliding and f is the function depending on the grain size distribution. Elongated grains or a non-uniform grain size distribution was often developed in the AI203 reported previously; the AI203 was not necessarily of such a high purity and not as easily sinterable as the present AI203 [11-14]. However, the high purity fine-grained AI203used in this study had a uniform grain shape and grain size distribution as shown in Fig. 1. No difference between the grain shapes and grain size distributions of the present A1203 and TZP reported previously was found. Consequently, it can be expected that the difference between the b values for AI203 and TZP results from c but not from f That is, the depth of deformation zone produced by unit sliding in TZP is about one twentieth of that in AI203. Such a difference may be related to the deformation temperature of the two materials, which is around 0.55 T m for TZP and about 0.70T m for AI203. It should be noted that the deformation temperature is also around 0.70 Tm for most superplastic metals. This conclusion is derived from the phenomenological analysis of strain-enhanced grain growth. However, the generation of a deformation zone due to grain boundary sliding has not yet been experimentally verified. Detailed microstructural examinations are required to make clear the mechanism of accelerated grain growth in ceramics. Finally, it must be pointed out that the model of Wilkinson and C~iceres is not fully applicable to strainenhanced grain growth in ceramics. For example, it is not possible to explain why the addition of 0.1 wt.% MgO causes a considerable decrease in enhanced grain growth rate. More important experimental evidence has been obtained in TZP by present authors. A small amount of glass addition to TZP changed the enhanced grain growth only slightly, despite a remarkable difference between the microstructures [23]. Grains in TZP were ordinarily faceted, but glass-containing specimens had round grains whose grain boundary junctions were filled with a glass phase. It is difficult to expect the generation of a deformation zone at the grain boundary corners on triple junctions in such a glass-containing TZR A full interpretation on the strain-enhanced grain
64
E Yoshizawa and T. Sakuma
/
Grain growth acceleration during high T deformation in Al_,O~
growth in ceramics has to be made by taking into account this experimental evidence.
5. Conclusion T h e stress-strain curves in high purity fine-grained Al20 3 was characterized by a high work-hardening rate after initial yielding, which was associated with grain growth during deformation. T h e grain growth was markedly accelerated during high temperature deformation, which was a linear function of plastic strain and was regarded as the strain-enhanced grain growth. T h e log-log plot of the strain-enhanced grain growth rate and the strain rate exhibited a linear relationship as in superplastic metals. T h e data in A l 2 0 3 were almost of the same magnitude as for superplastic metals, in contrast with T Z P which had a much lower strainenhanced grain growth rate. T h e difference between high purity A120 3 and T Z P was considered to be the difference in the grain boundary migration distance per unit grain boundary sliding.
Acknowledgments T h e authors wish to thank Mr. T. Seki for his experimental assistance. T h e authors also express their gratitude for the financial support by a Grant-in-Aid for Developmental Scientific Research ( 2)-01850154, for Scientific Research for Priority Areas (2)-02229206 and one of the authors (Y.Y.) for E n c o u r a g e m e n t of Young Scientists (A)-02750545 for Fundamental Scientific Research from the Ministry of Education, Science and Culture, Japan.
References 1 Y. Maehara and T. G. Langdon, J. Mater. Sci., 25 (1990) 2275.
2 l.-Wei Chen and L. A. Xue, J. Am. Ceram. Soc., 73 (1990) 2585. 3 F. Wakai, S. Sakaguchi and Y. Matsuno, Adv. Ceram. Mater., 1 (1986) 259. 4 B. J. Kellett and E E Lange, Z Am. Ceram. Soc., 69 (1986) C172. 5 T. Hermansson et al., in C. H. Hamilton and N. E. Paton (eds.), Superplasticity and Superplastic Forming, Minerals, Metals and Materials Society, Warrendale, PA, 1988, p. 24. 6 D. Dimos and D. L. Kohlstedt, J. Am. Ceram. Soc., 70 (1987) 531. 7 Y. Yoshizawa and T. Sakuma, Proc. 1st Jpn. Int. SAMPE Syrup., Nikkan Kogyo Shinbun, Tokyo, 1989, p. 272. 8 T. G. Nieh and J. Wadsworth, J. Am. Ceram. Soc., 72 (1989) 1469. 9 T. G. Nieh and J. Wadsworth, Acta Metall. Mater., 38 (1990) li21. 10 Y. Yoshizawa and T. Sakuma, Eng. Fract. Mech., (1991) in press. 11 R. M. Cannon, W. H. Rhodes and A. H. Heuer, J. Am. Ceram. Soc., 63 (1980) 46. 12 C. Carry and A. Mocellin, Ceram. Int., 13 (1987) 89. 13 J. D. Fridez, C. Carry and A. Mocellin, in W. D. Kingery (ed.), Structure and Properties of MgO and A1203 Ceramics, American Ceramic Society, Columbus, OH, 1984, p. 720. 14 K. R. Venkatachari and R. Raj, J. Am. Ceram. Soc., 69 (1986) 135. 15 E Wakai and H. Kato, Adv. Ceram. Mater., 3(1988) 71. 16 L. A. Xue and I.-Wei Chen, J. Am. Ceram. Soc., 73 (1990) 3518. 17 B. R Kashyap and K. Tangri, Metall. Trans. A, 18 (1987) 417. 18 C. H. C~iceres and D. S. Wilkinson, Acta Metall., 32 (1984) 415. 19 M.A. Clark and T. H. Alden, Acta Metall., 21 (1973) 1195. 20 O. N. Senkov and M. M. Myshlyaev, Acta Metall., 34 (1986) 97. 21 D. S. Wilkinson and C. H. Cficeres, Acta Metall., 32 (1984) 1335. 22 E. Sato, K. Kuribayashi and R. Horiuchi, in C. H. Hamilton and N. E. Paton (eds.), Superplasticity and Superplastic Forming, Mineral, Metals and Materials Society, Warrendale, PA, 1988, p. 115. 23 Y. Yoshizawa and T. Sakuma, Proc. Int. Conf. on Super° plasticity in Advanced Materials, 1991, 1991, in press.
59
Grain growth acceleration during high temperature deformation in high purity alumina Yu-ichi Yoshizawa and Taketo Sakuma Department of Materials Science, Faculty of Engineering, The Universityof Tokyo, 7-3-I Hongo, Tokyo 113 (Japan) (Received July 9, 1991 )
Abstract The accelerated grain growth during high temperature deformation in high purity alumina was investigated. The high purity alumina with an initial grain size about 0.6/~m was characterized by a high work-hardening rate after initial yielding at temperatures of 1250-1500 °C. The work hardening was related to the grain growth during deformation. The grain growth was accelerated by deformation which was analyzed to be of strain-enhanced type, because the grain size increase could be described as a function of strain but was not sensitive to temperature. The strain-enhanced grain growth in high purity AI203 and MgO-doped A1203 was discussed in comparison with the data for a tetragonal zirconia polycrystal reported previously.
1. Introduction Superplasticity in fine-grained ceramics, particularly in tetragonal zirconia polycrystals (TZPs), has been studied extensively by many workers [1-10]. It is essential to fabricate ceramics with a grain size below about 1/~m for realizing superplasticity in ceramics [1]. In the case of alumina (AI203), however, superplastic deformation has not been attained even for finegrained samples, because it is not easy to obtain steady state deformation owing to extensive grain growth and cavitation during high temperature deformation [11-16]. Such a microstructure change during deformation makes it difficult to understand the deformation mechanism of this material. In this study, the grain growth during high temperature deformation was examined in high purity A1203 and magnesia (MgO)doped AI203 in detail.
2. Experimental procedures High purity AI203 powders with an average diameter of 0.1 /xm supplied by Taimei Chemical Co. Ltd. (TM-DAR) and high purity MgO powders of 99.97% purity with about 17 nm diameter supplied by Ube Co. Ltd. were used as starting materials. Two kinds of material were prepared from these starting materials. One was pure A I 2 0 3 and the other was A I 2 0 3 doped with 0.1 wt.% MgO. AI203 powders were stirred with ethanol and followed by drying and passing through a 0921-5093/91/$3.50
60 mesh sieve. For fabricating MgO-doped A1203, the A1203 and MgO powders were mixed in a ball mill in ethanol together with high purity A120 3 balls for 24 h and then dried and sieved. These sieved powders were pressed into bars in a cemented carbide die under a pressure of 33 MPa and cold isostatically pressed under 100 MPa in a rubber tube. Sintering was done at 1300 °C for 2 h in air. The bulk density of the assintered and deformed specimens was measured by the Archimedes technique. Specimens of size about 7 m m x 7 m m x 7 mm were cut from the sintered bodies and ground with a diamond wheel. A high temperature mechanical test was done by uniaxial compression in a temperature range between 1250 and 1500 °C in air using an Instron-type Shimadzu AG5000C mechanical testing machine equipped with a high temperature furnace. The test was performed in a strain rate range between 1 × 10- 5 and 5 x 10- 4 s- 1. The compression test was started after holding at the test temperature for 10 rain. Small specimens, which were not deformed, were placed beside the compression test specimens to measure the static grain growth. The Pt-(Pt-Rh) thermocouples were attached to both specimens. The strain rate sensitivity exponent m was measured by the strain rate change method. Deformed and statically annealed specimens were cut and polished and were then thermally etched at 1200 °C for 4 h in air. They were coated with a gold film using an ion-sputtering machine in order to carry out scanning electron microscopy observations. The grain size was measured by the linear intercept method using © 1991 --Elsevier Sequoia, Lausanne
60
Y. Yoshizawa and T. Sakuma
/
Graingrowth acceleration during high T deformation in Al_,O~
photographs taken with a JEOL JSM-5200 scanning electron microscope. Metallographic examinations of thin foils prepared by ion milling were made with a JEOL JEM-2000FX transmission electron microscope operated at 200 kV.
3. Results
The density of the as-sintered state was about 98.5% of the theoretical value for pure and MgO-doped A1203, and the grain sizes of the two materials were 0.85/~m and 0.66/~m respectively. Figure 1 shows the transmission electron micrograph of MgO-doped A1203 sintered at 1300 °C for 2 h. The grains are nearly equiaxed and only a small number of pores is observed. No second phases such as glass phase or spinel were found, even at grain boundary triple junctions. The stress-strain curves for pure and MgO-doped A1203 deformed at 1300 °C for an initial strain rate of 1.3 x 10 -4 S- l are shown in Fig. 2. For comparison, the stress-strain curve of a TZP with an initial grain size of about 0.3/~m under the same deformation conditions is shown as a broken curve. The initial yield stresses (the 0.2% proof stress) of pure AI203 and MgO-doped A1203 are not very different from that of TZP at this temperature, but the work-hardening behaviours of A I 2 0 3 with and without MgO are significantly different from that of TZE A marked stress increase after initial yielding occurs in Al203 in contrast with TZP. The MgO addition results in a decrease in the work-hardening rate of A I 2 0 3. Figure 3 shows the stress-strain curves of MgOdoped Al203 at various temperatures. The compression test was done up to a maximum load of 20 kN.
The flow stress is largely dependent on the deformation temperature. The MgO-doped A1203 exhibits work hardening at temperatures between 1250 and 1500 °C. On the contrary, pure A1203 showed a larger work hardening than MgO-doped A I 2 0 3 as mentioned before. This large work hardening in pure AI203 often resulted in crack initiation, which is accompanied by a stress drop. Figure 4 is the transmission electron micrograph of deformed MgO-doped AI203. Grains remained equiaxed and dislocations are not observed. However, the grain size in the deformed sample is a little larger than in the as-sintered sample (Fig. 1). This microstructure was characteristic of fine-grained samples deformed at temperatures of 1250-1350 °C. In these samples, the density increased after deformation, e.g. MgO-doped AI203 with an initial density of 98.5% had densities of 99.7% and 99.9% after 30% strain at 1300 °C and 1350 °C respectively. Deformation above 1400 °C resulted in a different microstructure. Cavities were generally formed and the density
200
I
.~
i
1300"C
a_ 150
~uC .j °C100 ° ~ 50
AI203+HgO
jz_P_
..............
. . . . . . . . . . . .
0'3
o
lo STRA ] N
/
Fig. 2. Stress-strain curves of pure and MgO-doped AI203 at 1300 °C: - - - , stress-strain curve for TZP under the same conditions.
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/
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loo
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STRA l N Fig. 1. Transmission electron mirograph of MgO-doped AI~O3 sintered at 1300 °C for 2 h.
60
80
,
/
,
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Fig. 3. Stress-strain curves of MgO-doped AI203 at various temperatures.
Y. Yoshizawa and 7'. Sakuma
/
Grain growth acceleration during high T de[onnation in A 120 ~
61
Fig. 4. Transmission electron micrograph of MgO-doped AI203 31% deformed at 1300 °C. decreased after deformation above 1400 °C. In some coarse grains of size larger than about 2/~m, dislocations were introduced. Microstructural evidence indicates that it is not easy to understand the high temperature deformation mechanism in the present materials. In this study, grain growth behaviour was investigated for a grain size of less than 2/~m. Figures 5(a), 5(b) and 5(c) are the scanning electron micrographs of as-sintered, annealed and deformed MgO-doped A1203 respectively. The annealing and deformation temperature was 1300°C. In the assintered state, the average grain size was 0.66 tim. The grain sizes of both the annealed and the deformed samples are larger than that of the as-sintered sample. However, as clearly seen in Figs. 5(b) and 5(c), the deformed sample in Fig. 5(c) has a larger grain size than the annealed sample in Fig. 5(b). The average grain sizes were 0.67 ,um and 0.95 ktm for annealed and deformed samples respectively. This means that the grain growth is accelerated during high temperature deformation. The grain size change in MgO-doped AI203 during annealing and deformation at 1300 and 1350 °C is shown in Fig. 6. The terms static and dynamic are used for the grain growth during annealing and deformation in this figure. At 1300 °C, the static grain growth is almost negligible, but the grain size clearly increases with increasing deformation time. The grain growth takes place more rapidly at 1350 °C. At this temperature, the grain growth is also accelerated by deformation. 4. Discussion
It is well known that the grain growth during superplastic deformation in metals takes place faster than
Fig. 5. Scanning electron micrographs of MgO-doped AI203 (a) in the as-sintered state, (b) statically annealed and (c) 31% deformed at 1300 °C. The microstructure of (b) is obtained for a sample placed besides the deformed sample whose microstructure is (c). The time to keep at 1300 °C in the annealed sample is exactly the same as the loaded time in the deformed sample.
during static annealing [17-22]. The accelerated grain growth caused by high temperature deformation in superplastic metals is described as a function of plastic strain and is called the "strain-enhanced grain growth" [17]. The grain growth acceleration during superplastic or high temperature deformation has also been reported in TZP and AI20 ~ [2, 8, 9, 12-16]. Nieh and Wadsworth [8, 9] have indicated that the cube d 3 - d . ~ of grain size increment was proportional to the defor-
62
Y. Yoshizawa and 7". Sakuma
/
Graingrowth acceleration during high Tdeformation in AI,_O.~
1.5
1.5 1300" C 1350" C DYNAMI C • • / $TAT C O E] ]/,,,/"
:ZL
•
=1.0 ' v/'/'/ "/
/ ILl S
a
,
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Z
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- - O - - - O. . . . .
0.5(~
' 10
O - - O. . . . . . . . . .
' ~ 20 30 TIME /
L 4-0 rnin
, 50
60
Fig. 6. Grain size change during annealing and deformation at 1300 and 1350 °C for MgO-doped A1203: o, D, data for annealed samples; e,-, data for the deformed samples at an initial strain rate of 1.2 x 10 -4 S- t. The horizontal axis shows the time for annealing or deformation. mation time in TZP, where do is the initial grain size; Chen and Xue [2] have shown that In(d/do) changes with strain in Ce-TZE They have also reported that ( d - d o ) / d o is expressed as a function of deformation time in high purity A1203 [16]. Thus various plots have been proposed to analyse the accelerated grain growth in ceramics, but no theoretical background has been established in these plots. The present data in Fig. 6 clearly show that the dynamic grain growth includes both static growth and accelerated growth due to deformation. The contribution of static growth becomes more important at higher deformation temperatures. It seems therefore reasonable that the grain size difference d - d a between deformed and annealed samples should be discussed rather than d - d 0 for analysing the accelerated grain growth. The present authors have recently reported that the value of ( d - d a ) / d a increases almost linearly with strain and is not sensitive to deformation temperature in TZP [10, 23]. However, the magnitude of strainenhanced grain growth rate in TZP was about one order of magnitude lower than in most superplastic metals [10]. For examining whether or not the grain growth acceleration in high purity Al203 is expressed as a function of plastic strain, the normalized grain size increment ( d - d a ) / d a is plotted against strain in Fig. 7. The data on ( d - d a ) / d a for both pure and MgO-doped Al203 can be expressed by straight lines passing through the origin. The slopes of the lines are slightly different for pure and MgO-doped AI203 but are not sensitive to temperature. The slopes are estimated to be 2.4 and 1.5 for pure and MgO-doped Al203 respectively. These values are in good agreement with the reported value of 2 for superplastic Sn-Bi alloy
08.;
0.2
0.,
0.a
0.8
,0
Fig. 7. The relationship between the normalized grain size difference and strain for pure and MgO-doped A I 2 0 3 . The strain rate is in the range between 1.2 x 10 -5 and 5 x 10 -4 s - 1.
[19], but about 20 times larger than the value for TZP of about 0.1 [10]. Therefore it seems valid that the enhancement of grain growth caused by deformation in A1203 is regarded as strain enhanced. In comparison with superplastic metals, the experimental data on the strain-enhanced grain growth in ceramics are very limited, so that it is not possible to give a firm interpretation for the accelerated grain growth. However, for elucidating the mechanism of strain-enhanced grain growth, it may be valuable to examine whether or not the data for AI203 can be treated by analogy to those for superplastic metals. Wilkinson and Cficeres [21] have clarified a general rule on the strain-enhanced grain growth in superplastic metals. They have found that the log-log plot of grain growth rate normalized by initial grain size against strain rate is expressed by a sigmoidal relationship for various superplastic metals and is divided into three regions. In particular, the data fall in a narrow band and are expressed as a linear relationship in the intermediate strain rate range. Figure 8 is such a plot for the present alloys. In Fig. 8, the dotted zones are the data for various superplastic metals [21] and TZP obtained by the present authors [23]. In the strain rate range in Fig. 8, the data on both high purity A1203 and MgO-doped A120 3 are represented by straight lines whose slopes are close to unity as are those for superplastic metals and TZP. However, the absolute value of growth rate in alumina is much higher than that in TZP and rather agrees with that in superplastic metals.* *In the plot of Wilkinson and Cficeres, the strain-enhanced grain growth was not estimated from d-da as in the present paper but from d - d0. If the data on superplastic metals are replotted as d - d a, the conclusion may be slightly different.
Y. Yoshizawa and 7'. Sakuma ''"1
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Grain growth acceleration during high T deformation in AI,O
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63
deformed mainly by grain boundary sliding at the temperatures and strain rates examined. Then, it is rather realistic to assume that the parameters b, the average migration distance per unit sliding, are different in different materials and that b is much smaller in TZP than in A1203 and metals. It has been shown that b consists of two parameters:
'i!~
METAL (WILKINSON) '~
Ai203+MgO I • 1300' C
o
,
/
TZP b=cf
10-5 iii[
10-s
i
;: iii lili I iiiii
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i
i
i
I iiii[
i
i
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~ / s -1 Fig. 8. The relation between normalized strain-enhanced grain growth rate and strain rate. The data for metals are those compiled by Wilkinson and Cficeres [21 ], and those for TZP are the reported values obtained by present authors. The strain-enhanced grain growth rate in A1203 can be described by the following equation:
[t,. = abdg
(1)
where a is the fraction of overall strain due to grain boundary sliding, b is the average migration distance per unit sliding and d is the grain size. Equation (1) is for an intermediate strain rate range in superplastic metals, and the theoretical basis has been given by Wilkinson and Cficeres [21 ]. This equation predicts the proportionality between strain-enhanced grain growth rate and strain rate, and a linear relationship with a slope of unity in the log dr vs. log g plot being consistent with the data in Fig. 8. This may support the fact that the strain-enhanced grain growth in A1203 takes place by a similar mechanism to that in superplastic metals. Next, the difference of the proportional constant a b d in materials is discussed. Of the three parameters, the grain sizes d themselves are different in materials. A typical grain size for superplastic metals is several micrometres and that of TZP is 0.3-0.5 #m, while the grain size of the present AI203is between 0.6 and 1.5 #m. Thus it is better to discuss the enhanced grain growth by a normalized form:
d~, abg d
(2)
The proportional constant a b in eqn. (2) is estimated to be 2.5 for pure A1203, 1.4 for MgO-doped A1203, about 0.13 for TZP, and around 1 for superplastic metals. It is unrealistic to assume that the difference is caused by a, the fraction of overall strain due to sliding, because all these materials are believed to be
(3)
for grain growth enhanced by grain boundary sliding [23], where c is the depth of deformation zone produced by unit sliding and f is the function depending on the grain size distribution. Elongated grains or a non-uniform grain size distribution was often developed in the AI203 reported previously; the AI203 was not necessarily of such a high purity and not as easily sinterable as the present AI203 [11-14]. However, the high purity fine-grained AI203used in this study had a uniform grain shape and grain size distribution as shown in Fig. 1. No difference between the grain shapes and grain size distributions of the present A1203 and TZP reported previously was found. Consequently, it can be expected that the difference between the b values for AI203 and TZP results from c but not from f That is, the depth of deformation zone produced by unit sliding in TZP is about one twentieth of that in AI203. Such a difference may be related to the deformation temperature of the two materials, which is around 0.55 T m for TZP and about 0.70T m for AI203. It should be noted that the deformation temperature is also around 0.70 Tm for most superplastic metals. This conclusion is derived from the phenomenological analysis of strain-enhanced grain growth. However, the generation of a deformation zone due to grain boundary sliding has not yet been experimentally verified. Detailed microstructural examinations are required to make clear the mechanism of accelerated grain growth in ceramics. Finally, it must be pointed out that the model of Wilkinson and C~iceres is not fully applicable to strainenhanced grain growth in ceramics. For example, it is not possible to explain why the addition of 0.1 wt.% MgO causes a considerable decrease in enhanced grain growth rate. More important experimental evidence has been obtained in TZP by present authors. A small amount of glass addition to TZP changed the enhanced grain growth only slightly, despite a remarkable difference between the microstructures [23]. Grains in TZP were ordinarily faceted, but glass-containing specimens had round grains whose grain boundary junctions were filled with a glass phase. It is difficult to expect the generation of a deformation zone at the grain boundary corners on triple junctions in such a glass-containing TZR A full interpretation on the strain-enhanced grain
64
E Yoshizawa and T. Sakuma
/
Grain growth acceleration during high T deformation in Al_,O~
growth in ceramics has to be made by taking into account this experimental evidence.
5. Conclusion T h e stress-strain curves in high purity fine-grained Al20 3 was characterized by a high work-hardening rate after initial yielding, which was associated with grain growth during deformation. T h e grain growth was markedly accelerated during high temperature deformation, which was a linear function of plastic strain and was regarded as the strain-enhanced grain growth. T h e log-log plot of the strain-enhanced grain growth rate and the strain rate exhibited a linear relationship as in superplastic metals. T h e data in A l 2 0 3 were almost of the same magnitude as for superplastic metals, in contrast with T Z P which had a much lower strainenhanced grain growth rate. T h e difference between high purity A120 3 and T Z P was considered to be the difference in the grain boundary migration distance per unit grain boundary sliding.
Acknowledgments T h e authors wish to thank Mr. T. Seki for his experimental assistance. T h e authors also express their gratitude for the financial support by a Grant-in-Aid for Developmental Scientific Research ( 2)-01850154, for Scientific Research for Priority Areas (2)-02229206 and one of the authors (Y.Y.) for E n c o u r a g e m e n t of Young Scientists (A)-02750545 for Fundamental Scientific Research from the Ministry of Education, Science and Culture, Japan.
References 1 Y. Maehara and T. G. Langdon, J. Mater. Sci., 25 (1990) 2275.
2 l.-Wei Chen and L. A. Xue, J. Am. Ceram. Soc., 73 (1990) 2585. 3 F. Wakai, S. Sakaguchi and Y. Matsuno, Adv. Ceram. Mater., 1 (1986) 259. 4 B. J. Kellett and E E Lange, Z Am. Ceram. Soc., 69 (1986) C172. 5 T. Hermansson et al., in C. H. Hamilton and N. E. Paton (eds.), Superplasticity and Superplastic Forming, Minerals, Metals and Materials Society, Warrendale, PA, 1988, p. 24. 6 D. Dimos and D. L. Kohlstedt, J. Am. Ceram. Soc., 70 (1987) 531. 7 Y. Yoshizawa and T. Sakuma, Proc. 1st Jpn. Int. SAMPE Syrup., Nikkan Kogyo Shinbun, Tokyo, 1989, p. 272. 8 T. G. Nieh and J. Wadsworth, J. Am. Ceram. Soc., 72 (1989) 1469. 9 T. G. Nieh and J. Wadsworth, Acta Metall. Mater., 38 (1990) li21. 10 Y. Yoshizawa and T. Sakuma, Eng. Fract. Mech., (1991) in press. 11 R. M. Cannon, W. H. Rhodes and A. H. Heuer, J. Am. Ceram. Soc., 63 (1980) 46. 12 C. Carry and A. Mocellin, Ceram. Int., 13 (1987) 89. 13 J. D. Fridez, C. Carry and A. Mocellin, in W. D. Kingery (ed.), Structure and Properties of MgO and A1203 Ceramics, American Ceramic Society, Columbus, OH, 1984, p. 720. 14 K. R. Venkatachari and R. Raj, J. Am. Ceram. Soc., 69 (1986) 135. 15 E Wakai and H. Kato, Adv. Ceram. Mater., 3(1988) 71. 16 L. A. Xue and I.-Wei Chen, J. Am. Ceram. Soc., 73 (1990) 3518. 17 B. R Kashyap and K. Tangri, Metall. Trans. A, 18 (1987) 417. 18 C. H. C~iceres and D. S. Wilkinson, Acta Metall., 32 (1984) 415. 19 M.A. Clark and T. H. Alden, Acta Metall., 21 (1973) 1195. 20 O. N. Senkov and M. M. Myshlyaev, Acta Metall., 34 (1986) 97. 21 D. S. Wilkinson and C. H. Cficeres, Acta Metall., 32 (1984) 1335. 22 E. Sato, K. Kuribayashi and R. Horiuchi, in C. H. Hamilton and N. E. Paton (eds.), Superplasticity and Superplastic Forming, Mineral, Metals and Materials Society, Warrendale, PA, 1988, p. 115. 23 Y. Yoshizawa and T. Sakuma, Proc. Int. Conf. on Super° plasticity in Advanced Materials, 1991, 1991, in press.